Steel sheet with excellent cold workability during forming and method for manufacturing the same

ABSTRACT

The present invention provides a steel sheet having an excellent cold workability during forming and a method for producing the same. The steel sheet of the present invention is characterized in that: (a) the ratio of the number of carbides at the ferrite grain boundary to the number of carbides in the ferrite grain exceeds 1, (b) the ferrite grain diameter is 5 μm or more and 50 μm or less, (c) the in-plane anisotropy |Δr| of the r value is 0.2 or less, (d) the Vickers hardness is 100 HV or more and 150 HV or less, (e) the random intensity ratio of the {311} &lt;011&gt; orientation at the ½-thickness portion of the steel sheet is 3.0 or less.

TECHNICAL FIELD

The present invention relates to a steel sheet with excellent coldworkability during forming and a method for manufacturing the sheet.

BACKGROUND ART

Automotive parts, knives, and other mechanical parts are manufacturedthrough working processes such as punching, bending, and pressing. Inthe working processes, improvement of workability is required for amaterial carbon steel sheet, in order to improve product quality andstability and/or cost reduction.

Generally, a carbon steel sheet is subjected to cold rolling andspheroidizing annealing, so as to produce a soft carbon steel sheet withexcellent workability made of ferrite and spheroidized carbide. Manytechnologies for improving the workability of carbon steel sheets havebeen proposed so far.

For example, Patent Document 1 discloses a high-carbon steel sheet forprecision punching and a method for producing the sheet, wherein thesheet comprises, in terms of % by mass, C: 0.15 to 0.90%, Si: 0.40% orless, Mn: 0.3 to 1.0%, P: 0.03% or less, total Al: 0.1% or less, Ti:0.01 to 0.05%, B: 0.0005 to 0.0050%, N: 0.01% or less, and Cr: 1.2% orless, has a structure in which carbides having an average carbide grainsize of 0.4 to 1.0 μm and a carbide spheroidization ratio of 80% or moreare dispersed in a ferrite matrix, and has a notched tensile elongationof 20% or more.

Patent Document 2 discloses a medium- to high-carbon steel sheet withexcellent workability and a method for producing the sheet, wherein thesheet comprises C: 0.3 to 1.3 wt %, Si: 1.0 wt % or less, Mn: 0.2 to 1.5wt %, P: 0.02 wt% or less, and S: 0.02 wt % or less, has a structure inwhich carbides are dispersed so that the relationship C_(GB)/C_(IG)≤0.8holds between the carbide number C_(GB) on the ferrite crystal grainboundary and the carbide number C_(IG) in the ferrite crystal grains,and has a cross-sectional hardness of 160 HV or less.

Patent Document 3 discloses a medium- to high-carbon steel sheet withexcellent workability, wherein the sheet comprises C: 0.30 to 1.00 wt %,Si: 1.0 wt % or less, Mn: 0.2 to 1.5 wt %, P: 0.02 wt % or less, and S:0.02 wt % or less, has a structure in which carbides are dispersed inferrite so that the relationship C_(GB)/C_(IG)≤0.8 holds between thecarbide number C_(GB) on the ferrite crystal grain boundary and thecarbide number C_(IG) in the ferrite crystal grains, and simultaneously90% or more of the total carbides are occupied by spheroidized carbideshaving a long axis/short axis of 2 or less.

Patent Documents 1 to 3 describe that the greater the proportion ofcarbides in ferrite grains, the more the workability is improved.

In addition, Patent Document 4 discloses a steel sheet having excellentFB workability, mold life, and cold formability after FB processing,wherein the sheet comprises C: 0.1 to 0.5 wt %, Si: 0.5 wt % or less,Mn: 0.2 to 1.5 wt %, P: 0.03 wt % or less, S: 0.02 wt % or less, has astructure based on ferrite and carbide, and the amount S_(gb) of thecarbide present on the ferrite grain boundary is 40% or more, the aboveS_(gb) being defined by S_(gb)={S_(on)/(S_(on)+S_(in))}×100 (whereinS_(on) is the total area occupied by the carbides present on the grainboundary among the carbides present per unit area and S_(in) is thetotal area occupied by the carbides present on the grain boundary amongthe carbides present per unit area).

However, in the technology described in Patent Document 1, annealing isperformed at a temperature of the A_(C1) point or higher for softeningin order to coarsen ferrite grain size and carbide. But when annealingis performed at a temperature of the A_(C1) point or higher,rod-like/plate-like carbides may precipitate during annealing. Thecarbides, even though capable of reducing hardness, deteriorateworkability, which is disadvantageous in terms of workability.

The technologies described in Patent Documents 2 and 3 consider that thedeterioration of workability is caused by the low carbidespheroidization ratio of carbides precipitated on the grain boundary,but do not take into account the problem of improving thespheroidization ratio of grain boundary carbides. Techniques describedin Patent Document 4 only specify the tissue factor, and Patent Document4 does not discuss the relationship between workability and mechanicalproperties.

The technology described in Patent Document 5 is an invention made byfocusing on the relationship between fine blanking workability and theamount of carbide present in ferrite grains and ferrite grain size.However, Patent Document 5 does not discuss what effect the aggregatestructure has on the plastic anisotropy.

Patent Document 6 discloses a hot-rolled steel sheet in which thedevelopment of an aggregate structure otherwise developed by rolling issuppressed and a method for manufacturing the sheet. However, PatentDocument 6 does not discuss the relationship between the aggregatestructure other than the aggregate structure developed by rolling andthe cold forgeability.

The technology described in Patent Document 7 is an invention made byconsidering that the hardness and the total elongation of a high-carbonhot-rolled steel sheet prior to quenching are greatly influenced by thecementite density in the ferrite grains. The hot-rolled steel sheetdescribed in Patent Document 7 is characterized in that it has amicrostructure composed of ferrite and cementite, said microstructurehaving a cementite density of 0.10 strips/μm² or less in the ferritegrains. However, Patent Document 7 does not discuss what effect theaggregate texture has on the plastic anisotropy.

The technology described in Patent Document 8 is an invention made byconsidering that the C_(eq) value is related not only to mechanicalproperties and weldability but also to the fatigue crack growth rate insteels having a fine structure. Patent Document 8 discloses that bylimiting the range of the C_(eq) value to a range of 0.28% to 0.65%, thefatigue resistance of the steel material is improved and simultaneouslyweldability is secured. However, Patent Document 8 does not discuss whateffect the aggregate texture has on the plastic anisotropy.

PRIOR ART DOCUMENTS Patent documents

[Patent Document 1] Japanese Patent No. 4465057

[Patent Document 2] Japanese Patent No. 4974285

[Patent Document 3] Japanese Patent No. 5197076

[Patent Document 4] Japanese Patent No. 5194454

[Patent Document 5] Japanese Unexamined Patent Publication No.2007-270331

[Patent Document 6] Japanese Unexamined Patent Publication No.2009-263718

[Patent Document 7] Japanese Unexamined Patent Publication No.2015-17294

[Patent Document 8] Japanese Unexamined Patent Publication No.2004-27355

SUMMARY OF THE INVENTION Problem to be Solved by the Invention

In view of the current state of the prior art, it is an object of thepresent invention to address the problem of improving the coldworkability of a steel sheet during forming, and to provide a steelsheet that has solved the problem and a method for manufacturing thesheet.

Means to Solve the Problem

The present inventors have conducted intensive and extensive studies onmethods for solving the above-mentioned problems. As a result, thepresent inventors have found that by controlling the dispersion state ofthe carbide in the structure of the steel sheet before cold workingthrough the optimization of the manufacturing conditions in the stepsfrom hot rolling to annealing, the carbide can be precipitated on theferrite boundary and simultaneously the aggregate structure in the hotrolled steel plate can be controlled, thereby leading to enhanced coldworkability.

Further, we have found after intensive and extensive research that it isdifficult to manufacture a steel sheet that satisfies theabove-mentioned conditions merely by devising hot rolling conditions andannealing conditions separately, and that it can be manufactured byoptimizing the above conditions in mutual cooperation in an integratedprocess of the hot rolling and annealing steps.

The present invention has been made based on the above findings, and thegist thereof lies in:

(1) A steel sheet having an excellent cold workability during forming,comprising, in terms of % by mass:

C: 0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 0.30 to 1.00%, P: 0.0001 to0.020%, S: 0.0001 to 0.010%, Al: 0.001 to 0.10%, and

a balance of Fe and inevitable impurities,wherein (a) a ratio of the number of carbides at a ferrite grainboundary relative to the number of carbides in the ferrite grain is morethan 1,wherein (b) a diameter of the ferrite grain is 5 μm or more and 50 μm orless,wherein (c) an in-plane anisotropy |Δr| of the r value standardizedaccording to JIS Z 2254 is 0.2 or less,wherein (d) a Vickers hardness of the steel sheet is 100 HV or more and150 HV or less, andwherein (e) a ratio of X-ray diffraction intensity of the {311} <011>orientation at the ½-thickness portion of the steel sheet relative tothe X-ray diffraction intensity obtained when a sample with a randomorientation distribution of crystal grains in the steel sheet issubjected to X-ray diffraction is 3.0 or less.

(2) The steel sheet with excellent cold workability during formingdescribed in the above (1) further comprising, in terms of % by mass,one or a plurality of:

N: 0.0001 to 0.010%, O: 0.0001 to 0.020%, Cr: 0.001 to 0.50%, Mo: 0.001to 0.10%, Nb: 0.001 to 0.10%, V: 0.001 to 0.10%, Cu: 0.001 to 0.10%, W:0.001 to 0.10%, Ta: 0.001 to 0.10%, Ni: 0.001 to 0.10%, Sn: 0.001 to0.050%, Sb: 0.001 to 0.050%, As: 0.001 to 0.050%, Mg: 0.0001 to 0.050%,Ca: 0.001 to 0.050%, Y: 0.001 to 0.050%, Zr: 0.001 to 0.050%, La: 0.001to 0.050%, and Ce: 0.001 to 0.050%.

(3) A method for producing a steel sheet with excellent cold workabilityduring forming according to the above (1) or (2), the method comprising:

subjecting a steel strip having an ingredient composition according toclaim 1 or 2 to hot rolling by heating, followed by completing thefinish hot rolling at a temperature range of 800° C. or higher and 900°C. or lower;

coiling the hot-rolled steel sheet at a temperature of 400° C. or higherand 550° C. or lower;

pickling the hot-rolled steel sheet, and then subjecting the hot-rolledsteel sheet to a two-step type annealing in which the hot-rolled steelsheet is retained in two temperature ranges,

wherein the two-step type annealing comprises

(i) subjecting the hot-rolled steel sheet to a first step annealingperformed by retaining said hot-rolled steel at a temperature range of650° C. or higher and 720° C. or lower for 3 hours or longer and 60hours or shorter, and then a second step annealing performed byretaining the hot-rolled steel at a temperature range of 725° C. orhigher and 790° C. or lower for 3 hours or longer and 50 hours orshorter, and thereafter

(ii) cooling the hot-rolled steel sheet to 650° C. or lower at a coolingrate of 1° C./hour or more and 30° C./hour or less.

(4) The method for producing a steel sheet described in the above (3),wherein the steel sheet has a cross-sectional shrinkage percentage of40% or more.

Effect of the Invention

According to the present invention, a steel sheet with excellent coldworkability during forming can be manufactured and provided.

Mode for Carrying Out the Invention

A steel sheet with excellent cold workability during forming accordingto the present invention (hereinafter may be referred to as “theinventive steel sheet”) comprises, in terms of % by mass:

C: 0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 0.30 to 1.00%, P: 0.0001 to0.020%, S: 0.0001 to 0.010%, Al: 0.001 to 0.10%, and

a balance of Fe and inevitable impurities,

the above sheet being characterized in that:

(a) the ratio of the number of carbides at a ferrite grain boundaryrelative to the number of carbides in the ferrite grain exceeds 1,(b) the ferrite grain diameter is 5 μm or more and 50 μm or less,(c) the in-plane anisotropy |Δr| of the r value standardized accordingto JIS Z 2254 is 0.2 or less,(d) the Vickers hardness is 100 HV or more and 150 HV or less, and(e) the ratio of X-ray diffraction intensity of the {311}<011>orientation at the ½-thickness portion of the steel sheet relativeto the X-ray diffraction intensity obtained when a sample with a randomorientation distribution of crystal grains in the steel sheet issubjected to X-ray diffraction is 3.0 or less.

The method (hereinafter may be referred to as “the inventive method”) ofthe present invention for producing a steel sheet with excellent coldworkability during forming is a method for producing the inventive steelsheet,

wherein a hot-rolled steel strip that has been obtained by subjecting asteel strip having an ingredient composition of the inventive steelsheet to hot rolling by heating, followed by completing the finish hotrolling at a temperature range of 800° C. or higher and 900° C. orlower, and by coiling the resulting hot-rolled steel sheet at atemperature of 400° C. or higher and 550° C. or lower is, afterpickling, subjected to two-step type annealing in which the sheet isretained in two temperature ranges, whereupon

(i) the hot-rolled steel sheet is subjected to a first step annealingperformed by retaining said hot-rolled steel at a temperature range of650° C. or higher and 720° C. or lower for 3 hours or longer and 60hours or shorter, and then subjected to a second step annealingperformed by retaining the hot-rolled steel at a temperature range of725° C. or higher and 790° C. or lower for 3 hours or longer and 50hours or shorter, and thereafter

(ii) the sheet is cooled down to 650° C. or lower at a cooling rate of1° C./hour or more and 30° C./hour or less.

Hereinafter, the inventive steel sheet and the inventive manufacturingmethod will be described.

First, the reasons for limiting the ingredient composition of theinventive steel sheet will be described. The percentage relating to theingredient composition means % by mass.

C: 0.10 to 0.40%

C is an element that forms carbide in steel, and is effective forstrengthening steel and refining ferrite grains. In order to prevent thesurface of the steel sheet from being textured by cold working andensure the aesthetic appearance of surface of cold forged parts, it isnecessary to suppress the coarsening of ferrite grain size. However,when its content is less than 0.10%, the volume fraction of the carbideis insufficient and the coarsening of carbides during annealing cannotbe suppressed. Therefore, C is set to 0.10% or more, and preferably0.12% or more.

On the other hand, when it exceeds 0.40%, the volume fraction of thecarbide increases, a large amount of cracks serving as fracture startingpoints are formed when a load is instantaneously applied, and thus theimpact resistance property decreases. Therefore, C is set to 0.40% orless, and preferably 0.38% or less.

Si: 0.01 to 0.30%

Si is an element that acts as a deoxidizing agent and also affects theform of the carbide. In order to reduce the number of carbides in theferrite grain and increase the number of carbides on the ferrite grainboundaries, it is necessary to generate an austenite phase duringannealing in the two-step type annealing, and, after transientlydissolving the carbides, to cool gradually to promote the precipitationof carbides at the ferrite grain boundaries.

In the inventive steel sheet, the amount of Si may preferably be assmall as possible. However, when it is reduced to less than 0.01%, themanufacturing cost increases. Therefore, Si is set to 0.01% or more.

On the other hand, when it exceeds 0.30%, the ductility of ferritelowers and breaking may easily occur during cold working, resulting inreduced cold workability. Therefore, Si is set to 0.30% or less, andpreferably 0.28% or less.

Mn: 0.30 to 1.00%

Mn is an element that controls the figuration of carbides in thetwo-step type annealing. When its content is less than 0.30%, it isdifficult to precipitate carbides at the ferrite grain boundaries inslow cooling after the second-step annealing. Therefore, Mn is set to0.30% or more, and preferably 0.33% or more.

On the other hand, when it exceeds 1.00%, the hardness of ferriteincreases and the cold workability deteriorates. Therefore, Mn is set to1.00% or less, and preferably 0.96% or less.

P: 0.0001 to 0.020%

P is an element that segregates at the ferrite grain boundaries andsuppresses the formation of grain boundary carbides. The amount of P maypreferably be as small as possible. However, when P is reduced to lessthan 0.0001% in the refining process, the refining cost may greatlyincrease. Therefore, it is set to 0.0001% or more, and preferably0.0013% or more.

On the other hand, when it exceeds 0.020%, the number percentage of thegrain boundary carbides decreases and the cold workability deteriorates.Therefore, P is set to 0.020% or less, and preferably 0.018% or less.

S: 0.0001 to 0.010%

S is an element that forms a non-metallic inclusion such as MnS. Since anon-metallic inclusion serves as the starting point for break generationduring cold forging, the amount of S may preferably be as small aspossible. However, when it is reduced to less than 0.0001%, the refiningcost greatly increases. Therefore, S is set to 0.0001% or more, andpreferably 0.0012% or more.

On the other hand, when it exceeds 0.010%, cold workabilitydeteriorates. Therefore, S is set to 0.010% or less, and preferably0.007% or less.

Al: 0.001 to 0.10%

Al is an element that acts as a deoxidizing agent for steel andstabilizes ferrite. When its content is less than 0.001%, a sufficientaddition effect cannot be obtained. Therefore, Al is set to 0.001% ormore, and preferably 0.004% or more.

On the other hand, when it exceeds 0.10%, the number percentage ofcarbides on the grain boundary decreases and the cold workabilitydeteriorates. Therefore, Al is set to 0.10% or less, and preferably0.08% or less.

In addition to the above elements, the inventive steel sheet may containone or a plurality of N: 0.0001 to 0.010%, 0: 0.0001 to 0.020%, Cr:0.001 to 0.50%, Mo: 0.001 to 0.10%, Nb: 0.001 to 0.10%, V: 0.001 to0.10%, Cu: 0.001 to 0.10%, W: 0.001 to 0.10%, Ta: 0.001 to 0.10%, Ni:0.001 to 0.10%, Sn: 0.001 to 0.050%, Sb: 0.001 to 0.050%, As: 0.001 to0.050%, Mg: 0.0001 to 0.050%, Ca: 0.001 to 0.050%, Y: 0.001 to 0.050%,Zr: 0.001 to 0.050%, La: 0.001 to 0.050%, and Ce: 0.001 to 0.050%, inorder to improve the properties of the inventive steel sheet.

N: 0.0001 to 0.010%

N is an element that, when present in large amounts, causes theembrittlement of ferrite. The amount of N may preferably be as small aspossible. However, when it is reduced to less than 0.0001%, the refiningcost greatly increases. Therefore, N should be 0.0001% or more, andpreferably 0.0006% or more. On the other hand, when it exceeds 0.010%,ferrite embrittles and the cold forgeability deteriorates. Therefore, Nshould be 0.010% or less, and preferably 0.007% or less.

O: 0.0001 to 0.020%

O is an element that, when present in large amounts, forms coarse oxidesin steel. The amount of O may preferably be as small as possible.However, when it is reduced to less than 0.0001%, the refining costincreases greatly. Therefore, O is set to 0.0001% or more, andpreferably 0.0011% or more. On the other hand, when it exceeds 0.020%,coarse oxides are formed in the steel, the oxides serving as thestarting point for break generation during cold working. Therefore, 0 isset to 0.020% or less, and preferably 0.017% or less.

Cr: 0.001 to 0.50%

Cr is an element which enhances quenchability and contributes to theimprovement of strength and which is thickened to carbide and formsstable carbide even in the austenitic phase. When its content is lessthan 0.001%, the sufficient effect of improving quenchability cannot beobtained. Therefore, Cr is set to 0.001% or more, and preferably 0.007%or more. On the other hand, when it exceeds 0.50%, the carbide becomesstabilized thereby delaying the dissolution of the carbide duringquenching, and thus, it is feared that the desired quenching strengthmay not be achieved. Therefore, Cr is set to 0.50% or less, andpreferably 0.45% or less.

Mo: 0.001 to 0.10%

Like Mn, Mo is an element effective for controlling the figuration ofcarbides. When its content is less than 0.001%, a sufficient additioneffect cannot be obtained. Therefore, Mo is set to 0.001% or more, andpreferably 0.010% or more. On the other hand, when it exceeds 0.10%, thein-plane anisotropy of the r value deteriorates and the cold workabilitydeteriorates. Therefore, Mo is set to 0.10% or less, and preferably0.08% or less.

Nb: 0.001 to 0.10%

Nb is an element which is effective for controlling the figuration ofcarbides and which refines the structure, thereby contributing to theenhancement of its toughness. When its content is less than 0.001%, asufficient addition effect cannot be obtained. Therefore, Nb should be0.001% or more, and preferably 0.004% or more. On the other hand, whenit exceeds 0.10%, a large number of fine Nb carbides precipitate, whichleads to excessively increased strength. It also causes the reduction inthe number ratio of grain boundary carbides, and the deterioration incold forgeability. Therefore, Nb is set to 0.10 or less, and preferably0.08% or less.

V: 0.001 to 0.10%

Like Nb, V is an element which is effective for controlling thefiguration of carbides and which refines the structure, therebycontributing to the enhancement of its toughness. When its content isless than 0.001%, a sufficient addition effect cannot be obtained.Therefore, V is set to 0.001% or more, and preferably 0.004% or more. Onthe other hand, when it exceeds 0.10%, a large number of fine V carbidesprecipitate, which leads to excessively increased strength, to thereduced number ratio of grain boundary carbides, and to the deterioratedcold forgeability. Therefore, V is set to 0.10 or less, and preferably0.08% or less.

Cu: 0.001 to 0.10%

Cu is an element which segregates at the ferrite crystal grain boundaryand forms fine precipitates thereby to contribute to the enhancement ofstrength. When its content is less than 0.001%, a sufficient effect ofenhancing strength cannot be obtained. Therefore, Cu is set to 0.001% ormore, and preferably 0.005% or more. On the other hand, when it exceeds0.10%, red heat embrittlement occurs and the productivity by hot rollingdecreases. Therefore, Cu is set to 0.10% or less, and preferably 0.08%or less.

W: 0.001 to 0.10%

Like Nb and V, W is also an element effective for controlling thefiguration of carbides. When its content is less than 0.001%, asufficient addition effect cannot be obtained. Therefore, W is set to0.001% or more, and preferably 0.003% or more. On the other hand, whenit exceeds 0.10%, a large number of fine W carbides precipitate, whichleads to excessively increased strength, to the reduced number ratio ofgrain boundary carbides, and to the deteriorated cold forgeability.Therefore, W is set to 0.10 or less, and preferably 0.08% or less.

Ta: 0.001 to 0.10%

Like Nb, V and W, Ta is also an element effective for controlling thefiguration of carbides. When its content is less than 0.001%, asufficient addition effect cannot be obtained. Therefore, W is set to0.001% or more, and preferably 0.005% or more. On the other hand, whenit exceeds 0.10%, a large number of fine W carbides precipitate, whichleads to excessively increased strength, to the reduced number ratio ofgrain boundary carbides, and to the deteriorated cold forgeability.Therefore, Ta is set to 0.10 or less, and preferably 0.08% or less.

Ni: 0.001 to 0.10%

Ni is an element effective for improving the toughness of parts. Whenits content is less than 0.001%, a sufficient addition effect cannot beobtained. Therefore, Ni is set to 0.001% or more, and preferably 0.003%or more. On the other hand, when it exceeds 0.10%, the number ratio ofgrain boundary carbides decreases and the cold forgeabilitydeteriorates. Therefore, Ni is set to 0.10% or less, and preferably0.08% or less.

Sn: 0.001 to 0.050%

Sn is an element contaminated from a steel raw material (scrap). Itsegregates at the grain boundary, leading to the decreased number ratioof grain boundary carbides. Therefore, its content may preferably be assmall as possible. However, when it is reduced to less than 0.001%, therefining cost will be greatly increased. Therefore, Sn is set to 0.001%or more, and preferably 0.002% or more. On the other hand, when itexceeds 0.050%, ferrite embrittles and cold forgeability deteriorates.Therefore, Sn is set to 0.050% or less, and preferably 0.040% or less.

Sb: 0.001 to 0.050%

Like Sb, Sb is an element contaminated from a steel raw material(scrap). It segregates at the grain boundary, leading to the decreasednumber ratio of grain boundary carbides. Therefore, its content maypreferably be as small as possible. However, when it is reduced to lessthan 0.001%, the refining cost will be greatly increased. Therefore, Sbis set to 0.001% or more, preferably 0.002% or more. On the other hand,when it exceeds 0.050%, the cold forgeability deteriorates. Therefore,Sb is set to 0.050% or less, and preferably 0.040% or less.

As: 0.001 to 0.050%

Like Sn and Sb, As is an element contaminated from a steel raw material(scrap). It segregates at the grain boundary, thereby leading to adecrease in the number ratio of grain boundary carbides. Therefore, itscontent may preferably be as small as possible. However, when it isreduced to less than 0.001%, the refining cost increases greatly.Therefore, As is set to 0.001% or more, and preferably 0.002% or more.On the other hand, when it exceeds 0.050%, the number ratio of the grainboundary carbides decreases and the cold forgeability deteriorates.Therefore, As is set to 0.050% or less, and preferably 0.040% or less.

Mg: 0.0001 to 0.050%

Mg is an element that can control the figuration of sulfides with theaddition of its trace amount. When its content is less than 0.0001%, asufficient addition effect cannot be obtained. Therefore, Mg is set to0.0001% or more, and preferably 0.0008% or more. On the other hand, whenit exceeds 0.050%, ferrite embrittles and the cold forgeabilitydeteriorates. Therefore, Mg is set to 0.050% or less, and preferably0.040% or less.

Ca: 0.001 to 0.050%

Like Mg, Ca is an element that can control the figuration of sulfideswith the addition of its trace amount. When its content is less than0.001%, a sufficient addition effect cannot be obtained. Therefore, Cais set to 0.001% or more, and preferably 0.003% or more. On the otherhand, when it exceeds 0.050%, coarse Ca oxides are formed, which serveas starting points of break generation during cold forging. Therefore,Ca is set to 0.050% or less, and preferably 0.040% or less.

Y: 0.001 to 0.050%

Like Mg and Ca, Y is an element that can control the figuration ofsulfides with the addition of its trace amount. When its content is lessthan 0.001%, a sufficient addition effect cannot be obtained. Therefore,Y is set to 0.001% or more, and preferably 0.003% or more. On the otherhand, when it exceeds 0.050%, coarse Y oxides are formed, which serve asstarting points of break generation during cold working. Therefore, Y isset to 0.050% or less, and preferably 0.035% or less.

Zr: 0.001 to 0.050%

Like Mg, Ca and Y, Zr is an element that can control the figuration ofsulfides with the addition of its trace amount. When its content is lessthan 0.001%, a sufficient addition effect cannot be obtained. Therefore,Zr is set to 0.001% or more, and preferably 0.004% or more. On the otherhand, when it exceeds 0.050%, coarse Zr oxides are formed, which serveas starting points for break generation during cold working. Therefore,Zr is set to 0.050% or less, and preferably 0.045% or less.

La: 0.001 to 0.050%

La is an element that can control the figuration of sulfides with theaddition of its trace amount, but it is also an element that segregatesat the grain boundary and causes a decrease in the number ratio of grainboundary carbides. When its content is less than 0.001%, a sufficienteffect of controlling figuration cannot be obtained. Therefore, La isset to 0.001% or more, and preferably 0.004% or more. On the other hand,when it exceeds 0.050%, the number ratio of grain boundary carbidesdecreases and the cold workability deteriorates. Therefore, La is set to0.050% or less, and preferably 0.045% or less.

Ce: 0.001 to 0.050%

Like La, Ce is an element that can control the figuration of sulfideswith the addition of its trace amount, but it is also an element thatsegregates at the grain boundary and causes a decrease in the numberratio of grain boundary carbides. When its content is less than 0.001%,a sufficient effect of controlling figuration cannot be obtained.Therefore, Ce is set to 0.001% or more, and preferably 0.004% or more.On the other hand, when it exceeds 0.050%, the number ratio of grainboundary carbides decreases and the cold forgeability deteriorates.Therefore, Ce is set to 0.050% or less, and preferably 0.045% or less.

The remainder of the ingredient composition of the inventive steel sheetis Fe and unavoidable impurities.

It is a novel finding by the inventors that the inventive steel sheethas excellent cold workability during forming, because, in addition tothe above ingredient composition, it was found, as a result of optimumhot rolling and annealing, that

(a) the ratio of the number of carbides at the ferrite grain boundaryrelative to the number of carbides in the ferrite grain exceeds 1,

(b) the ferrite grain diameter is 5 μm or more and 50 μm or less,

(c) the in-plane anisotropy |Δr| of the r value standardized accordingto JIS Z 2254 is 0.2 or less,

(d) the Vickers hardness is 100 HV or more and 150 HV or less, and

(e) the ratio of X-ray diffraction intensity of the {311} <011>orientation at the ½-thickness portion of the steel sheet relative tothe X-ray diffraction intensity obtained when a sample with a randomorientation distribution of crystal grains in the steel sheet issubjected to X-ray diffraction is 3.0 or less.

The above (a) to (e) will be described below.

(a) The ratio of the number of carbides at the ferrite grain boundaryrelative to the number of carbides in the ferrite grain exceeds 1:

The inventive steel sheet has a structure which is substantiallycomposed of ferrite and carbide, and in which the ratio of the number ofcarbides at the ferrite grain boundary relative to the number ofcarbides in the ferrite grain exceeds 1. Carbides are, in addition tocementite (Fe₃C) that is a compound of iron and carbon, compoundsobtained by replacing Fe in cementite with an element such as Mn and Cr,and alloy carbides (M₂₃C₆, M₆Co, MC, etc., wherein M is Fe and anotheradditive metal element).

When a steel sheet is formed into a predetermined part shape, a shearband is formed in the macrostructure of the steel sheet, and slipdeformation is generated and concentrated in the vicinity of the shearband. The slip deformation involves propagation of dislocations, andregions with high dislocation density are formed in the vicinity of theshear band. As the strain amount applied to the steel sheet increases,the slip deformation is promoted and thereby the dislocation densityincreases. In cold forging, strong processing exceeding an equivalentstrain of 1 is applied.

For this reason, in the conventional steel sheet, generation of voidsand/or cracks due to the increased dislocation density could not beprevented, and it was difficult to improve cold forgeability.

In order to solve the above challenging problems, it is effective tosuppress the formation of shear bands during forming. From the viewpointof a microstructure, shear band formation is a phenomenon in which aslip generated in one crystal grain crosses the crystal grain boundaryand propagates continuously to an adjacent crystal grain. Therefore, inorder to suppress the formation of a shear band, it is necessary toprevent the propagation of slippage beyond crystal the grain boundary.

Carbides in the steel sheet are tenacious particles that hinderslippage. Therefore, the presence of carbides at the ferrite grainboundaries would make it possible, for the first time, to suppress theformation of a shear band and thereby to improve cold forgeability.

Based on the theory and principle, it is considered that coldforgeability is strongly influenced by the coverage rate of carbides atthe ferrite grain boundaries. Therefore, it becomes necessary to measurethe coverage rate with high accuracy.

In order to measure the coverage rate of carbides at the ferrite grainboundaries in a three-dimensional space, serial sectioning SEMobservation or repeated three-dimensional EBSP observation becomeessential in which sample cutting by FIB and observation are repeated inthe scanning electron microscope. However, these methods take a hugeamount of measurement time and the accumulation of technical know-howbecomes indispensable. We clarified this fact and concluded that commonanalytical methods are not suitable.

Therefore, as a result of searching a simple and highly accurateevaluation index, the present inventors have found that coldforgeability can be evaluated by using, as an index, the ratio of thenumber of carbides at the ferrite grain boundary relative to the numberof carbides in the ferrite grain, and that cold forgeability can beremarkably improved when the ratio of the number of carbides at theferrite grain boundary relative to the number of carbides in the ferritegrain is more than 1.

Any of buckling, folding and convolution of a steel sheet that occursduring cold working is caused by the localization of strain accompanyingthe formation of a shear band. Therefore, by allowing the carbide toexist at the ferrite grain boundaries, the formation of the shear bandand the localization of strain can be alleviated and the generation ofbuckling, folding and convolution can be suppressed.

When the spheroidization percentage of carbides on the crystal grainboundary is less than 80%, strains are concentrated locally on therod-shaped or plate-shaped carbides, and voids and/or cracks are likelyto occur. Therefore, the carbide spheroidization ratio on the crystalgrain boundary may preferably be 80% or more, and more preferably 90% ormore.

When the average particle diameter of the carbide in the ferrite grainand the carbide at the ferrite grain boundaries is less than 0.1 μm, thehardness of the steel sheet remarkably increases and the workabilitydeteriorates. Therefore, the average particle diameter of the carbidemay preferably be 0.1 μm or more, and more preferably 0.17 μm or more.On the other hand, when the average particle diameter of the carbideexceeds 2.0 μm, fissures occur with the coarse carbide serving as astarting point during cold working, and thus the cold workabilitydeteriorates. Therefore, the average particle diameter of the carbidemay preferably be 2.0 μm or less, and more preferably 1.95 μm or less.

Subsequently, the method of observing and measuring the structure willbe described.

Observation of the carbide is carried out by a scanning electronmicroscope. Prior to observation, samples for structure observation arepolished by wet polishing with emery paper and polishing with diamondabrasive grains having an average particle size of 1 μm. After polishingthe observation surface to a mirror finish, the structure is etched witha 3% nitric acid-alcohol solution.

Among the magnification for observation, within 3000 times, amagnification capable of discriminating between ferrite and carbide isselected. At the selected magnification, eight images with a viewingfield of 30 μm×40 μm are randomly photographed at the ¼ plate layerthickness.

With respect to the tissue image obtained, the area of each carbidecontained in the region is measured in detail by an image analysissoftware represented by Mitsuya Shoji Co. Ltd. (Win ROOF). A circleequivalent diameter (=2×√(area/3.14)) is obtained from the area of eachcarbide, and the average value is taken as the carbide particlediameter.

Further, the spheroidization ratio of the carbide was determined byapproximating the carbide to an ellipse having an equal area and equalmoment of inertia, and then by calculating the proportion of thecarbides in which the ratio of the maximum length to the maximum lengthin the perpendicular direction is less than 3.

In order to suppress the effect of measurement error due to noise,carbides having an area of 0.01 μm² or more among the carbides in grainsand grain boundaries were counted and the carbides having an area of0.01 μm² or less were excluded from evaluation.

The number of carbides present on the ferrite grain boundary wascounted, and from the total number of carbides the number of carbides inthe ferrite grain was determined by subtracting the number of carbideson the ferrite grain boundary. Based on the measured number, the ratioof the number of carbides on the grain boundary relative to the numberof carbides in the ferrite grain was determined.

(b) The ferrite grain diameter is 5 μm or more and 50 μm or less:

In the structure after annealing the cold rolled steel sheet, the coldworkability can be improved by setting the ferrite grain diameter to 5μm or more. When the ferrite grain size is less than 5 μm, the hardnessincreases and fissures and cracks tend to generate easily during coldworking. Therefore, the ferrite grain size is set to 5 μm or more, andpreferably 7 μm or more.

On the other hand, when it exceeds 50 μm, the number of carbides on thecrystal grain boundary that suppress slippage propagation decreases andthe cold workability deteriorates. Therefore, that the ferrite grainsize is set to 50 μm or less, and preferably 37 μm or less.

The ferrite grain diameter is measured in the above-described polishingmethod, wherein the observation surface of the sample is polished to amirror surface, followed by etching with a 3% nitric acid-alcoholsolution. The structure of the observation surface is then examined withan optical microscope or a scanning electron microscope, and a linesegment method is then applied to the image photographed to determinethe ferrite grain diameter.

(c) The in-plane anisotropy |Δr| of the r value standardized accordingto JIS Z 2254 is 0.2 or less:

The in-plane anisotropy |Δr| of the plastic strain ratio (r value) ofthe steel sheet is measured in a method in accordance with JIS Z 2254.The r value (0° direction: r₀, 45° direction: r₄₅, 90° direction: r₉₀)measured by taking test strips from each direction of 0° direction, 45°direction and 90° direction with respect to the rolling direction wasused to calculate the following equation.

|Δr|=(r ₀−2r ₄₅ +r ₉₀)/2

By setting the in-plane anisotropy |Δr| of the plastic strain ratio (rvalue) of the steel sheet to 0.2 or less, the cold workability can beimproved. When |Δr| exceeds 0.2, the thickness of parts and the heightof the earing become uneven during drawing. Therefore, the in-planeanisotropy |Δr| is set to 0.2 or less.

(d) The Vickers hardness is 100 HV or more and 150 HV or less:

By setting the Vickers hardness of the steel sheet to 100 HV or more and150 HV or less, the cold workability can be improved. When the Vickershardness is less than 100 HV, buckling can easily occur during coldworking. Therefore, the Vickers hardness is set to 100 HV or more, andpreferably 110 HV or more.

On the other hand, when the Vickers hardness exceeds 150 HV, theductility decreases and the internal breaking tends to occur easilyduring cold forging. Therefore, the Vickers hardness is set to 150 HV orless, and preferably 146 HV or less.

(e) The ratio of X-ray diffraction intensity of the {311} <011>orientation at the ½-thickness portion of the steel sheet relative tothe X-ray diffraction intensity obtained when a sample with a randomorientation distribution of crystal grains in the steel sheet issubjected to X-ray diffraction is 3.0 or less:

In cold forging, in addition to controlling the figuration of carbides,the draw formability during cold forging must be secured. In order toimprove the draw formability during cold forging, plastic anisotropysuch as in-plane anisotropy |Δr| must be improved. For that purpose, theaggregate structure of a hot-rolled steel sheet must be controlled. Forevaluation of the aggregate structure, analysis by X-ray diffraction ona plane parallel to the plate surface at the ½ thickness portion of thehot-rolled steel plate is used.

One surface of a hot-rolled steel plate is ground to a ½ plate thicknesssurface in parallel to the surface to expose a ½ plate thicknesssurface, followed by the analysis of the ½ plate thickness surface byX-ray diffraction. As the X-ray diffraction, X-ray diffraction by Mobulb may be used. Diffraction intensities of diffraction orientations{110}, {220}, {211} and {310} by reflection are obtained, and basedthereon, the orientation distribution function (ODF) is created.

The X-ray diffraction intensity ratio is determined by using thediffraction intensity data of the ½ plate thickness surface obtainedfrom the ODF and the diffraction intensity data of random orientation ofthe hot-rolled steel sheet. Specifically, as a standard sample in whichthe metallic structure has no accumulation in a specific direction, asample obtained by sintering powder iron of a hot-rolled steel sheet tobe measured or the powder before sintering is used to determine thediffraction intensity under the same conditions as when the diffractionintensity data of the ½ plate thickness surface was obtained. The partto be collected as the standard sample is not particularly limited andmay be any part of the hot-rolled steel sheet. The X-ray diffractionintensity ratio in a specific orientation is a numerical value obtainedby dividing the diffraction intensity in the specific direction of the ½plate thickness surface obtained from the ODF by the diffractionintensity of the standard sample.

When the X-ray diffraction intensity ratio of the {311} <011>orientation obtained by the above-described ODF analysis is set to I1,it is necessary that this I1 is 3.0 or less, and preferably 2.5 or lessfor the random aggregate structure during hot rolling. When a randomaggregate structure having I1 of 3.0 or less can be obtained, theplasticity anisotropy is reduced and the cold formability is improved.

Next, the inventive manufacturing method will be described.

The manufacturing method according to the present invention ischaracterized in that the hot rolling and the annealing are consistentlymanaged to control the structure. After continuously casting a steelstrip having a predetermined ingredient composition, the steel strip issubjected to hot rolling by heating to complete finish hot rolling at atemperature range of 800° C. or higher to 900° C. or lower, coiled at400° C. or higher and 550° C. or lower to obtain a hot-rolled steelsheet. The hot-rolled steel sheet is, after pickling, subjected to atwo-step type annealing in which the hot-rolled steel sheet ismaintained in two temperature ranges, whereupon

(i) the hot-rolled steel sheet is subjected to a first step annealingperformed by retaining said hot-rolled steel at a temperature range of650° C. or higher and 720° C. or lower for 3 hours or longer and 60hours or shorter, and then subjected to a second step annealingperformed by retaining the hot-rolled steel at a temperature range of725° C. or higher and 790° C. or lower for 3 hours or longer and 50hours or shorter, and thereafter

(ii) the hot-rolled steel sheet is cooled to 650° C. or lower at acooling rate of 1° C./hour or more and 30° C./hour or less,

and thus a steel sheet excellent in cold workability during forming canbe produced.

By the hot rolling and annealing mentioned above, a structure composedof fine pearlite and bainite can be formed as the structure of the steelsheet.

The processing conditions will be described below.

Heating temperature of a steel strip: 1000° C. or higher and 1250° C. orlower

The heating temperature of the steel strip subjected to hot rolling maypreferably be 1000° C. or higher and 1250° C. or lower, and the heatingtime may preferably be 0.5 hour or longer and 3 hours or shorter.

When the heating temperature is lower than 1000° C. or the heating timeis shorter than 0.5 hour, the microsegregation and/or macrosegregationformed by casting are not eliminated, and regions in which Si, Mn, etc.,are locally concentrated inside the steel material may remain, and thusthe impact resistance property of the steel material is lowered.Therefore, the heating temperature may preferably be 1000° C. or higher,and preferably 0.5 hour or longer.

On the other hand, when the heating temperature exceeds 1250° C. or theheating time exceeds 3 hours, decarburization from the surface layer ofthe steel strip becomes conspicuous, and austenite grains in the surfacelayer grow abnormally during heating before carburizing and quenching,and the impact resistance property of the steel strip is deteriorated.Thus the heating temperature may preferably be 1250° C. or lower, andthe heating time may preferably be 3 hours or shorter.

Finish hot rolling temperature: 800° C. or higher and 900° C. or lower

Finish hot rolling is completed at 800° C. or higher and 900° C. orlower. When the finish hot rolling temperature is lower than 800° C.,the deformation resistance of the steel strip increases, the rollingload increases markedly, the wear amount of the roll increases, and theproductivity decreases. Therefore, the finish hot rolling temperature isset to 800° C. or higher, and preferably 820° C. or higher.

On the other hand, when the finish hot rolling temperature exceeds 900°C., thick scales are generated during plate passing on the ROT (Run OutTable), scratches are generated on the surface of the steel sheet due tothe scale, and cracks are generated starting from scratches when animpact load is applied after cold forging and carburizing and annealing,leading to reduced impact resistance property of the steel sheet.Therefore, the finish hot rolling temperature is set to 900° C. orlower, and preferably 880° C. or lower.

Cooling rate on ROT: 10° C./sec or more and 100° C./sec or less

The cooling rate at the time of cooling the hot-rolled steel sheet onthe ROT after finish hot rolling may preferably be 10° C./sec or moreand 100° C./sec or less. When the cooling rate is less than 10° C./sec,thick scales are generated during cooling and the occurrence ofscratches on the surface of the steel sheet due to the scales cannot besuppressed. Therefore, the cooling rate is set to 10° C./sec or more,and more preferably 20° C./sec or more.

On the other hand, when the cooling rate exceeds 100° C./sec, the steelsheet is cooled at a cooling rate exceeding 100° C./sec from the surfacelayer to the inside of the steel sheet, the outermost layer part of thesteel sheet is excessively cooled, and a low-temperature transformedstructure such as bainite or martensite is formed.

At the time of discharging the hot-rolled coil cooled from 100° C. toroom temperature after coiling, microcracks are generated in thelow-temperature transformed structure. It is difficult to remove themicrocracks in the subsequent pickling step and cold rolling step, andfissures progress from the microcracks as a starting point during coldworking, leading to reduced cold workability. Therefore, the coolingrate may preferably be 100° C./sec or less.

Note that the above cooling rate refers to the cooling capacity from thecooling facility at each water injection zone from the point at whichthe hot-rolled steel sheet after the finish hot rolling is cooled at thewater injection zone after passing through the water-free zone to apoint at which it is cooled to the coiling target temperature on theROT, and does not refer to the average cooling rate from the waterinjection starting point to the temperature at which it is coiled by thecoiling device.

Coiling temperature: 400° C. or higher and 550° C. or lower

The coiling temperature is set to 400° C. or higher and 550° C. orlower. When the coiling temperature is lower than 400° C., the austenitewhich was not transformed before coiling is transformed into hardmartensite, cracks are generated in the surface layer of the steel sheetduring discharge of the hot-rolled coil, leading to reduced workability.Therefore, the coiling temperature is set to 400° C. or higher, andpreferably 430° C. or higher.

On the other hand, when the coiling temperature exceeds 550° C.,pearlite having a large lamellar spacing is generated and thickneedle-shaped carbides having high thermal stability are formed, andeven after the two-step type annealing, needle-shaped carbides remain.Since fissures are generated during cold working with theseneedle-shaped carbides as a starting point, the coiling temperature isset to 550° C. or lower, and preferably 520° C. or lower.

The hot-rolled coil manufactured under the above conditions is annealed,after pickling, in a two-step type annealing which retains the coil intwo temperature ranges. The first-step annealing and the second-stepannealing may be either box annealing or continuous annealing. Bycontrolling the stability of carbides by the two-step type annealing,the formation of carbides on the ferrite grain boundary and thespheroidization ratio of carbides on the ferrite grain boundary can beenhanced.

The two-step type annealing will be described below.

The first step annealing is carried out in a temperature range of theA_(ci) point or lower to coarsen carbides and enrich alloy elements toincrease the thermal stability of carbides. Thereafter, the temperatureis raised to a range from A_(C1) point or higher to A₃ point or lower togenerate austenite in the structure.

Thereafter, by gradual cooling, the austenite is transformed intoferrite and the carbon concentration in the austenite is increased. Byproceeding slow cooling, carbon atoms are adsorbed to the carbidesremaining in the austenite, and thus the carbide and austenite come tocover the grain boundary of the ferrite. Finally it becomes possible toform a structure in which many spheroidized carbides are present in thegrain boundary of the ferrite.

When the residual carbides are small in quantity while maintaining thetemperature range of A_(C1) point or higher to A₃ point or lower,pearlite, rod-shaped carbides and plate-like carbides are producedduring cooling. When these pearlite, rod-shaped carbides and plate-likecarbides are produced, the workability of the steel sheet is remarkablydeteriorated. Therefore, increasing the number of residual carbides inthe temperature range from A_(C1) point or higher to A₃ point or loweris an important factor to enhance the workability of the steel sheet.

By using a steel sheet structure obtained under the above hot rollingcondition, the thermal stability of carbides at a temperature of A_(C1)point or lower can be secured. Therefore, an increase in the number ofresidual carbides in the temperature range from A_(C1) point or higherto A₃ point or lower can be targeted.

Hereinafter, an annealing condition for the two-step type annealing willbe described.

First step annealing

Temperature range: 650° C. or higher and 720° C. or lower

Retention time: 3 hours or longer and 60 hours or shorter

In the first step annealing, the annealing temperature is set to 650° C.or higher and 720° C. or lower. When the annealing temperature of thefirst step is lower than 650° C., the stability of the carbide becomesinsufficient and it becomes difficult to allow the carbide to remain inthe austenite in the second step annealing. Therefore, the temperatureof the first step annealing is set to 650° C. or higher, and preferably670° C. or higher.

On the other hand, when the temperature of the first step annealingexceeds 720° C., austenite is generated before enhancing the stabilityof the carbide, which makes it difficult to control the required changein the structure. Therefore, the first step annealing temperature is setto 720° C. or lower, and preferably 700° C. or lower.

The retention time at the first step is 3 hours or longer and 60 hoursor shorter. When the retention time is lower than 3 hours, the stabilityof the carbide is insufficient and it becomes difficult to allow thecarbide to remain at the second step annealing. Therefore, the retentiontime of the first step is set to 3 hours or longer. On the other hand,when the retention time of the first step exceeds 60 hours, improvementof the stability of the carbide cannot be expected and furthermore theproductivity is lowered. Therefore, the retention time of the first stepis set to 60 hours or shorter, and preferably 55 hours or shorter.

The annealing atmosphere is not limited to a specific atmosphere. Forexample, it may be either a nitrogen atmosphere having a nitrogencontent of 95% or more, a hydrogen atmosphere having a hydrogen contentof 95% or more, or an atmospheric atmosphere.

Second step annealing

Temperature range: 725° C. or higher and 790° C. or lower

Retention time: 3 hours or longer and 50 hours or shorter

In the second step annealing, the annealing temperature is set to 725°C. or higher and 790° C. or lower. When the second-step annealingtemperature is lower than 725° C., the amount of austenite produced issmall and the number ratio of carbides on the ferrite grain boundary islowered. Therefore, the second-step annealing temperature is set to 725°C. or higher, and preferably 745° C. or higher.

On the other hand, when the second-step annealing temperature exceeds790° C., it becomes difficult to allow the carbide to remain in theaustenite and to control the required structure change. Therefore, thesecond-step annealing temperature is set to 790° C. or lower, andpreferably 770° C. or lower.

The retention time of the second step is set to 3 hours or longer and 50hours or shorter. When the retention time of the second step is lessthan 3 hours, the amount of austenite produced is small, dissolution ofthe carbide in the ferrite grains is insufficient, and it becomesdifficult to increase the number ratio of carbides on the ferrite grainboundary. Therefore, the retention time of the second step is set to 3hours or longer, and preferably 5 hours or longer.

On the other hand, when the retention time of the second step exceeds 50hours, it becomes difficult to allow the carbide to remain in theaustenite. Therefore, the retention time of the second step is set to 50hours or shorter, and preferably is 46 hours or shorter.

The annealing atmosphere is not limited to a specific atmosphere. Forexample, it may be either a nitrogen atmosphere having a nitrogencontent of 95% or more, a hydrogen atmosphere having a hydrogen contentof 95% or more, or an atmospheric atmosphere.

After completion of the two-step type annealing, the hot-rolled steelsheet is cooled, whereupon it is cooled to 650° C. at a cooling rate of1° C./hour or more to 30° C./hour or less.

Cooling rate to a temperature of 650° C. or lower: 1° C./hour or moreand 30° C./hour or less

Since the temperature range for controlling the structure change by slowcooling is sufficient up to 650° C., it is only necessary to control thecooling rate in the temperature range up to 650° C. After reaching atemperature of 650° C. or lower, it may be cooled to room temperaturewithin the above range without controlling the cooling rate.

It may be preferable that the cooling rate is slow in order to graduallycool the austenite produced in the second step annealing to transforminto ferrite and allow carbon to be adsorbed to the carbides remainingin the austenite. However, when the cooling rate is less than 1°C./hour, the time required for cooling increases and the productivitydecreases. Therefore, the cooling rate is 1° C./hour or more, andpreferably 5° C./hour.

On the other hand, when the cooling rate exceeds 30° C./hour, austenitetransforms to pearlite, the hardness of the steel sheet increases, thecold forgeability deteriorates, and the impact resistance property ofthe steel sheet after carburizing quenching and tempering decreases.Therefore, the cooling rate is set to 30° C./hour or less, andpreferably 26° C./hour or less.

Further, according to the inventive production method, a steel sheetwith excellent cold workability during forming can be produced in whichthe ingredient composition is, in terms of % by mass, comprising: C:0.10 to 0.40%, Si: 0.01 to 0.30%, Mn: 0.30 to 1.00% , P: 0.0001 to0.020%, S: 0.0001 to 0.010%, and Al: 0.001 to 0.10%, the balance beingFe and unavoidable impurities, the metal structure is substantiallycomposed of ferrite and spheroidized carbides, and (a) the ratio of thenumber of carbides at the ferrite grain boundary to the number ofcarbides in the ferrite grain exceeds 1, (b) the ferrite grain size is5μm or more and 50 μm or less, (c) the in-plane anisotropy |Δr| of the rvalue standardized according to JIS Z 2254 is 0.2 or less, (d) theVickers hardness is 100 HV or more and 150 HV or less, thecross-sectional shrinkage percentage is 40% or more, and the ratio ofX-ray diffraction intensity of the {131} <011> orientation at the½-thickness portion of the steel sheet relative to the X-ray diffractionintensity obtained when a sample having the random orientationdistribution of crystal grains in the steel sheet is subjected to X-raydiffraction is 3.0 or less.

The cross-sectional shrinkage percentage is defined by the followingformula (1). A large value of this ratio means that the localdeformability is high, and as the value of the formula (1) increases,the workability of the steel sheet increases.

Sectional shrinkage percentage (%)=100-(cross-sectional area at tensilefracture/initial cross-sectional area)×100   Equation (1)

As described above, the present invention is characterized in that byrolling control and heat treatment after rolling, a structure in whichcarbides (that is, cementite) are uniformly dispersed is formed, so thatthe crystal anisotropy can be eliminated. Therefore, in the presentinvention, the random intensity ratio of the {311} <011> orientation atthe ½ plate thickness portion of the steel sheet can be made 3.0 orless.

EXAMPLES

Next, examples will be described, but the level of examples is anexample of conditions adopted for confirming the feasibility andeffectiveness of the present invention, and the present invention is notlimited to this one condition example. The present invention can adoptvarious conditions as long as the object of the present invention isachieved without departing from the gist of the present invention.

Example 1

In order to investigate the effect of hot rolling conditions, acontinuous cast strip (steel ingot) having the ingredient compositionshown in Table 1 was subjected to hot rolling under the conditions shownin Table 2 to produce a hot-rolled coil having a thickness of 3.0 mm.Incidentally, the steel type described as “Developed steel” in thecolumn of “Remarks” in Table 1 has a composition included in thecomposition range of the steel sheet according to the present invention.Also, the steel type described as “Comparative steel” in the column of“Remarks” in Table 1 has a composition outside the composition range ofthe steel sheet according to the present invention. In addition, theingredients that do not satisfy the composition conditions of the steelsheet according to the present invention are underlined.

A sample for characterization was prepared as follows: a hot-rolledcoil, after pickling, was placed in a box-type annealing furnace, theatmosphere was controlled to 95% hydrogen-5% nitrogen, the coil washeated from room temperature to 705° C. and was retained for 36 hours tomake the temperature distribution uniform in the hot-rolled coil. Thecoil was then heated to 760° C. and retained at 760° C. for 10 hours,and was then cooled to 650° C. at a cooling rate of 10° C./hour, thenfurnace-cooled to room temperature to prepare the sample forcharacterization. The structure of the sample was measured by the methoddescribed above.

TABLE 1 C Si Mn P S Al N O Cr Mo Nb V Cu W A 0.21 0.13 0.53 0.00480.0084 0.618 0.0053 0.098 0.084 B 0.24 0.28 0.91 0.0039 0.0018 C 0.340.19 0.51 0.0133 0.0015 0.049 D 0.18 0.18 0.50 0.0151 0.0092 0.00810.461 E 0.32 0.03 0.73 0.0149 0.0094 0.018 0.0018 0.087 0.080 F 0.220.06 0.98 0.0144 0.0088 0.082 0.073 0.039 G 0.17 0.13 0.94 0.0199 0.00260.038 0.277 H 0.37 0.11 0.38 0.0144 0.0087 I 0.33 0.10 0.96 0.00100.0019 1.180 0.480 J 0.29 0.16 0.81 0.0132 0.0031 K 0.11 0.02 0.410.0117 0.0360 L 0.25 0.11 0.63 0.0167 0.0058 0.0007 0.090 0.020 M 0.191.21 0.91 0.0027 0.0084 N 0.31 0.23 1.27 0.0010 0.0019 O 0.51 0.14 0.650.0059 0.0013 P 0.20 0.13 0.92 0.0153 0.0091 Q 0.11 0.07 0.64 0.00600.0051 0.0013 0.0131 R 0.27 0.19 0.70 0.0185 0.0033 0.012 S 0.27 0.130.80 0.0044 0.0002 0.071 0.060 T 0.25 0.14 0.38 0.0142 0.0097 U 0.120.25 0.86 0.0154 0.0045 W 0.40 0.08 0.78 0.0184 0.0087 0.026 0.025 0.0880.055 X 0.13 0.20 0.96 0.0061 0.0034 Y 0.21 0.21 0.41 0.0139 0.00380.092 0.093 Z 0.11 0.05 0.59 0.0195 0.0010 AA 0.13 0.09 0.76 0.01270.0079 AB 0.19 0.22 0.89 0.0005 0.0021 AC 0.12 0.02 0.63 0.0168 0.0098AD 0.30 0.28 0.71 0.0169 0.0031 0.018 0.035 Ta Ni Sn Sb As Mg Ca Y Zr LaCe Remarks A 0.004 0.015 0.0149 Comparative steel B Developed steel C0.083 0.018 0.0459 Developed steel D 0.017 0.018 Developed steel E 0.0200.034 0.021 0.045 Developed steel F Developed steel G 0.014 Developedsteel H Developed steel I 0.037 0.046 0.048 Comparative steel JDeveloped steel K Comparative steel L 0.012 0.0423 0.035 0.008 Developedsteel M 0.037 0.046 0.048 Comparative steel N Comparative steel OComparative steel P Developed steel Q 0.033 0.035 Developed steel R0.080 0.037 Developed steel S 0.016 0.0280 0.036 Developed steel TDeveloped steel U Developed steel W 0.038 0.015 Developed steel XDeveloped steel Y 0.039 0.030 Developed steel Z Developed steel AADeveloped steel AB Developed steel AC Developed steel AD 0.041 Developedsteel

TABLE 2 Hot rolling Condition Grain Cross- {311}<011> Finish hot Ferriteboundary sectional X-ray rolling Coiling Carbide grain Vickers carbideshrinkage diffraction In-plane temp. temp. diameter diameter hardnessNo./grain percent intensity anisotropy [° C.] [° C.] [μm] [μm] [HV]carbide No. [%] ratio (I1) |Δr| Remarks A-1 891 505 0.91 18.5 125 3.751.0 5.5 0.44 Comparative steel B-1 832 500 1.17 20.7 124 6.7 49.7 1.60.05 Inventive steel C-1 850 489 1.08 20.7 120 4.6 51.3 2.1 0.10Inventive steel D-1 742 532 0.99 21.5 111 6.3 54.0 5.6 0.48 Comparativesteel E-1 811 617 0.95 16.8 116 5.4 29.7 1.5 0.06 Comparative steel F-1937 512 1.18 21.1 112 7.2 52.8 2.5 0.18 Comparative steel G-1 871 4091.27 23.9 110 8.7 56.3 1.5 0.04 Inventive steel H-1 837 522 0.99 20.7117 3.7 55.7 2.0 0.10 Inventive steel I-1 802 484 1.03 16.9 121 7.1 49.55.0 0.39 Comparative steel J-1 841 407 1.28 22.5 116 6.9 43.9 1.9 0.09Inventive steel K-1 871 475 1.11 29.7 104 4.4 31.4 1.4 0.06 Comparativesteel L-1 898 443 1.20 22.7 110 6.9 54.7 2.3 0.12 Inventive steel M-1873 467 1.11 18.1 172 5.4 34.9 4.6 0.37 Comparative steel N-1 850 4141.37 22.8 152 9.3 38.5 2.4 0.14 Comparative steel O-1 834 413 1.22 20.0127 5.7 38.4 2.0 0.12 Comparative steel P-1 856 485 1.22 22.8 111 7.257.1 2.2 0.11 Inventive steel Q-1 873 428 1.25 29.3 116 6.2 54.1 2.00.12 Inventive steel R-1 845 348 1.32 24.0 114 6.5 51.1 2.2 0.14Comparative steel S-1 837 454 1.23 21.9 114 7.1 56.6 1.8 0.08 Inventivesteel T-1 830 671 0.75 17.4 116 2.8 38.7 1.4 0.05 Comparative steel U-1881 419 1.29 27.6 112 7.3 59.0 2.7 0.19 Inventive steel W-1 853 535 1.0918.5 121 6.1 55.1 1.4 0.05 Inventive steel X-1 885 477 1.24 25.8 112 7.558.0 2.9 0.19 Inventive steel Y-1 874 436 1.14 25.7 110 5.8 44.8 2.00.12 Inventive steel Z-1 899 485 1.16 27.7 108 5.5 57.5 1.6 0.06Inventive steel AA-1 895 423 1.28 27.6 102 6.9 42.6 2.4 0.14 Inventivesteel AB-1 855 434 1.27 23.9 115 7.3 53.3 2.5 0.16 Inventive steel AC-1895 381 1.30 30.1 109 6.4 54.4 1.9 0.11 Comparative steel AD-1 854 4631.17 20.7 125 5.9 52.3 1.4 0.06 Inventive steel

The cold workability was evaluated using the notched tensile test andthe in-plane anisotropy of the r value. In the notched tensile test, anotched tensile test strip was taken from an as-annealed material with athickness of 3 mm, and a tensile test was performed in the rollingdirection to determine the cross-sectional shrinkage percentage, and thelocal deformability was evaluated. When the cross-sectional shrinkagepercentage is 40% or more, it was rated as superior.

Further, the in-plane anisotropy of the r value was rated as superiorwhen the in-plane anisotropy |Δr| of the r value standardized accordingto JIS Z 2254 of an as-annealed material with a thickness of 3 mm was0.2 or less.

In order to determine the X-ray diffraction intensity ratio (I1) of{311} <011>, X-ray diffraction with an Mo tube was performed from thecenter of the plate thickness of each sample followed by an ODFanalysis. Based on the results obtained by the ODF analysis, the I1 wasdetermined.

Table 2 shows, for each of the samples prepared, the results of thecarbide diameter, the ferrite grain diameter, the Vickers hardness, theratio of the number of carbides at the ferrite grain boundary relativeto the number of carbides in the ferrite grain, the cross-sectionalshrinkage percentage, the X-ray diffraction intensity ratio of {311}<011> and in-plane anisotropy. Among the samples in Table 2, thoseindicated as “Inventive steel” in the Remarks column satisfy therequirements of the steel sheet according to the present invention, andthose indicated as “Comparative steel” in the Remarks column do notsatisfy the requirements of the steel sheet according to the presentinvention. In Table 2, the measurement results that do not satisfy therequirements of the steel sheet according to the present invention andthe manufacturing conditions that do not satisfy the requirements of thesteel sheet manufacturing method according to the present invention areunderlined.

As shown in Table 2, in any of the inventive steels B-1, C-1, G-1, H-1,J-1, L-1, P-1, Q-1, S-1, U-1, W-1, X-1, Y-1, Z-1, AA-1, AB-1 and AD-1,the ratio of the number of carbides at the ferrite grain boundaryrelative to the number of carbides in the ferrite grain exceeds 1, andthe Vickers hardness is 150 HV or less. In addition, in any of theinventive steels, the cross-sectional shrinkage percentage exceeds 40%and the in-plane anisotropy |Δr| of the r value is 0.2 or less. Thus,they have excellent cold workability. Furthermore, since it wasconfirmed that scale scratches were not generated on the steel sheetsurface in any of the inventive steels, these steels can be suitablyused for cold working.

On the other hand, in the Comparative steel A-1, since the Al content ishigh and the A3 point decreased, recrystallization during finish hotrolling was inhibited and |Δr| deteriorated. Thus, the cold workabilityis low. In the Comparative steel I-1, the contents of Mo and Cr arehigh, recrystallization during finish hot rolling was inhibited, and|Δr| deteriorated. In the comparative steels K-1 and N-1, the content ofS or Mn is high, coarse MnS was formed in the steel, and the coldworkability is low. In the Comparative steel M-1, the content of Si washigh and hardness increased, and thus cold workability is low. Also, inthe Comparative steel M-1, since the A3 point rose, recrystallizationduring finish hot rolling was hindered and |Δr| deteriorated.

In the Comparative steel O-1, C is high, the volume fraction of carbidesincreased, a large amount of cracks as the starting point of fractureswere generated, and the cross-sectional shrinkage percentage was low.Thus, the cold workability is low. In the Comparative steel D-1, thefinish temperature of hot rolling was low and the productivitydecreased. In the Comparative steel F-1, the finish temperature of hotrolling was high, and scale scratches were generated on the surface ofthe steel sheet.

In the Comparative steels R-1 and AC-1, the coiling temperature of hotrolling was low, the low-temperature transformation structure such asbainite and martensite increased resulting in brittled steel, and breaksfrequently occurred when the hot-rolled coil was discharged resulting ina decrease in productivity. In the Comparative steels E-1 and T-1, thecoiling temperature of hot rolling was high, thick pearlite withlamellar spacing and needle-shaped coarse carbides with high thermalstability were produced in the hot rolled structure. Since thesecarbides remained in the steel sheet even after the two-step typeannealing, the cross-sectional shrinkage percentage was low and thus thecold workability is low.

Subsequently, in order to investigate the effect of annealingconditions, steel strips (slabs) having the ingredient composition shownin Table 1 were heated at 1240° C. for 1.8 hours and then subjected tohot rolling. After completing finish hot rolling at 890° C., they werecooled to 520° C. at a cooling rate of 45° C./sec on ROT and coiled at510° C. to produce a hot-rolled coil with a thickness of 3.0 mm. Andunder the conditions shown in Table 3, a hot-rolled sheet-annealedsample with a thickness of 3.0 mm was prepared.

For each of the samples prepared, the carbide diameter, the ferritegrain diameter, the Vickers hardness, the ratio of the number ofcarbides at the ferrite grain boundary relative to the number ofcarbides in the ferrite grain, the cross-sectional shrinkage percentage,the X-ray diffraction intensity ratio of {311} <011> and the in-planeanisotropy were determined in the same manner as the inventive steelsand the comparative steels in Table 2. The results are shown in Table 3.

TABLE 3 1st step annealing 2nd step annealing Ferrite RetentionRetention Retention Retention Cooling Carbide grain temp. time temp.time rate diameter diameter [° C.] [hr] [° C.] [hr] [° C./sec] [μm] [μm]A-2 669 27 761 41  7 0.96 27.1 B-2 695 47 753 12 30 0.80 15.2 C-2 654 19771 32 20 0.90 23.3 D-2 705 51 760 23 13 0.98 25.4 E-2 693 50 729 45 340.60 13.6 F-2 698 22 743 36 14 0.91 16.8 G-2 695 55 759 19  7 1.26 24.2H-2 698 30 742  1  6 0.98 13.5 I-2 658 57 754 14 16 0.75 12.8 J-2 694 42811 29  7 1.59 38.1 K-2 675 21 746 42 22 0.71 26.5 L-2 694 14 781 12 300.84 22.0 M-2 709  9 779  8 29 0.81 18.2 N-2 670 15 738 14 25 0.76 10.9O-2 676 52 732 46  7 0.82 14.6 P-2 671 21 769 42  9 1.23 29.5 Q-2 701  2756 34 25 0.79 24.9 R-2 663 52 729 15 17 0.61 10.3 S-2 681 56 779  4  71.41 27.3 T-2 741 47 774  9 25 1.29 32.6 U-2 672 12 741 47 25 0.71 17.7W-2 700 45 730  8 28 0.72 10.3 X-2 709 46 743 45  8 1.15 22.2 Y-2 676 51777 54 17 0.94 37.7 Z-2 668 30 706 23 16 0.57  8.5 AA-2 662 14 776 26 192.23 54.0 AB-2 678 68 785 12 14 1.16 29.4 AC-2 637 35 745 48  6 1.1926.2 AD-2 712 50 750 29 15 0.96 19.3 Grain {311}<011> boundary Cross-X-ray Vickers carbide sectional diffraction In-plane hardness No./grainshrinkage intensity anisotropy [HV] carbide No. ratio [%] ratio (I1)|Δr| Remarks A-2 116 4.8 56.3 4.6 0.38 Comp. steel B-2 136 3.1 48.7 1.90.10 Developed steel C-2 112 2.8 58.7 1.8 0.11 Developed steel D-2 1083.7 59.4 1.5 0.06 Developed steel E-2 161 1.7 38.6 1.2 0.04 Comp. steelF-2 120 4.9 54.8 2.5 0.17 Developed steel G-2 118 11.4  51.0 2.2 0.15Developed steel H-2 133 0.4 50.9 1.5 0.07 Comp. steel I-2 133 8.2 53.63.8 0.28 Comp. steel J-2 151 2.9 32.6 1.9 0.10 Comp. steel K-2 104 2.720.7 1.0 0.03 Comp. steel L-2 109 4.3 55.9 2.4 0.17 Developed steel M-2183 3.9 38.6 2.8 0.18 Comp. steel N-2 153 11.2  39.6 1.6 0.06Comp. steel O-2 148 4.0 37.4 1.2 0.05 Comp. steel P-2 108 4.6 59.0 1.40.05 Developed steel Q-2 135 3.2 37.6 1.6 0.05 Comp. steel R-2 138 4.548.0 2.0 0.12 Developed steel S-2 126 7.6 49.9 1.2 0.04 Developed steelT-2 138 1.5 34.7 1.1 0.01 Comp. steel U-2 121 3.7 42.1 2.6 0.16Developed steel W-2 141 2.1 46.5 2.0 0.11 Developed steel X-2 118 4.249.6 1.8 0.08 Developed steel Y-2 141 2.8 29.2 1.8 0.07 Comp. steel Z-2124 0.5 51.0 1.1 0.04 Comp. steel AA-2 105 9.4 60.2 1.4 0.05 Developedsteel AB-2 116 4.8 45.7 1.4 0.06 Comp. steel AC-2 133 6.8 31.2 1.9 0.10Comp. steel AD-2 127 2.3 50.9 1.7 0.06 Developed steel

As shown in Table 3, in any of the inventive steels B-2, C-2, D-2, F-2,G-2, L-2, P-2, R-2, S-2, U-2, W-2, X-2, AA-2 and AD-2, the ratio of thenumber of carbides at the ferrite grain boundary to the number ofcarbides in the ferrite grain exceeds 1, and the Vickers hardness is 150HV or less. In addition, in any of the inventive steels, thecross-sectional shrinkage percentage exceeds 40% and the in-planeanisotropy |Δr| of the r value is 0.2 or less. Thus, they have excellentcold workability.

On the other hand, in the Comparative steel A-2, since the Al content ishigh and the A3 point decreased, recrystallization during finish hotrolling was inhibited and |Δr| deteriorated. Thus, the cold workabilityis low. In the Comparative steel 1-2, the contents of Mo and Cr arehigh, recrystallization during finish hot rolling was inhibited, and|Δr| deteriorated. In the comparative steels K-2 and N-2, the content ofS or Mn is high, coarse MnS was formed in the steel. Thus, the coldworkability deteriorated. In the Comparative steel M-2, the content ofSi was high and hardness increased. Thus, the cold workability is low.Also, in the Comparative steel M-2, since the A3 point decreased,recrystallization during finish hot rolling was hindered and |Δr|deteriorated.

In the Comparative steel O-2, C is high, the volume fraction of carbidesincreased, a large amount of cracks as the starting point of fracturewere generated, and the cross-sectional shrinkage percentage was low.Thus, the cold workability is low.

In the Comparative steel AC-2, since the annealing temperature in thefirst-step annealing during the two-step type box annealing is low, thetreatment of carbide coarsening at the Ac1 temperature or lower isinsufficient, and the thermal stability of carbides is insufficient,thus the carbides remaining at the second step of annealing decreases,the pearlite transformation cannot be suppressed in the structure afterthe slow cooling, and the cross-sectional shrinkage percentage is low.Thus, the cold workability is low.

In the Comparative steel T-2, since the annealing temperature in thefirst step annealing during the two-step type box annealing is high,austenite is generated during annealing and the stability of carbidecannot be increased, so that carbides remaining during the second stepannealing decrease, and pearlite transformation cannot be suppressed inthe structure after the slow cooling, and the cross-sectional shrinkagepercentage is low. Thus, the cold forgeability is low.

In the Comparative steel Q-2, since the retention time in the first stepannealing during annealing of the two-step type is short, the treatmentof the carbide coarsening at the Ac1 temperature or lower isinsufficient, and the thermal stability of the carbide is insufficient,and thus the carbide remaining at the second step of annealing decreasesand the pearlite transformation cannot be suppressed in the structureafter the slow cooling, and the cross-sectional shrinkage percentage islow. And thus the cold workability is low. In the comparative steelAB-2, the retention time during the first stage box annealing of thetwo-step type is long and the productivity is low.

In the Comparative steel Z-2, since the annealing temperature during thesecond-step annealing during the two-step type box annealing is low, andthe amount of austenite produced is small, so that the proportion of thenumber of carbides in the grain boundary cannot be increased. Thus, thecold workability is low. In the Comparative steel J-2, the annealingtemperature during the second-step annealing during the two-step typeannealing is high, the amount of the carbide remaining is decreased dueto the promoted dissolution of carbides, and pearlite transformationcannot be suppressed in the structure after the slow cooling, theVickers hardness is too high, and the cross-sectional shrinkagepercentage is low. Thus, the cold forgeability is low.

In the Comparative steel H-2, since the annealing temperature during thesecond-step annealing during the two-step type annealing is low, and theamount of austenite produced is small, so that the proportion of thenumber of carbides in the grain boundary cannot be increased. Thus, thecold workability is low. In the Comparative steel Y-2, the retentiontime during the second-step annealing during the two-step type annealingis long, the amount of carbides remaining is decreased due to thepromoted dissolution of carbides, and pearlite transformation cannot besuppressed in the structure after the slow cooling, and thecross-sectional shrinkage percentage is low. Thus, the cold forgeabilityis low. In the Comparative steel E-2, the cooling rate from thesecond-step annealing during the two-step type annealing to 650° C. isfast, pearlite transformation occurred during cooling, the Vickershardness is too high, and the cross-sectional shrinkage percentage islow. Thus, the cold workability is low.

In any of the comparative steels A-1, D-1, I-1, M-1, A-2 and I-2, theX-ray diffraction intensity ratio of {311} <011> is greater than 3.0. Inthese comparative steels, the in-plane anisotropy |Δr| exceeds 0.2, andthus the cold workability is low. As described above, by performinganalysis by X-ray diffraction on a plane parallel to the plate surfaceat the ½ plate thickness portion of the hot-rolled steel sheet, thedegree of plastic anisotropy such as the in-plane anisotropy |Δr| or thequality of cold workability of the hot-rolled steel sheet to be coldworked can be determined before cold working.

INDUSTRIAL APPLICABILITY

As described above, according to the present invention, a steel sheetwith excellent cold workability during forming can be manufactured andprovided. The steel sheet of the present invention is a steel sheetsuitable as a material for automotive parts, blades, and othermechanical parts manufactured through processing steps such as punching,bending, pressing, etc. Therefore, the present invention has excellentindustrial applicability.

1. A steel sheet having an excellent cold workability during forming,comprising, in terms of % by mass: C: 0.10 to 0.40%, Si: 0.01 to 0.30%,Mn: 0.30 to 1.00%, P: 0.0001 to 0.020%, S: 0.0001 to 0.010%, Al: 0.001to 0.10%, and a balance of Fe and inevitable impurities, wherein (a) aratio of the number of carbides at a ferrite grain boundary relative tothe number of carbides in the ferrite grain is more than 1, wherein (b)a diameter of the ferrite grain is 5 μm or more and 50 μm or less,wherein (c) an in-plane anisotropy |Δr| of the r value standardizedaccording to JIS Z 2254 is 0.2 or less, wherein (d) a Vickers hardnessof the steel sheet is 100 HV or more and 150 HV or less, and wherein (e)a ratio of X-ray diffraction intensity of the {311} <011> orientation atthe ½-thickness portion of the steel sheet relative to the X-raydiffraction intensity obtained when a sample with a random orientationdistribution of crystal grains in the steel sheet is subjected to X-raydiffraction is 3.0 or less.
 2. The steel sheet with excellent coldworkability during forming according to claim 1 further comprising, interms of % by mass, one or a plurality of: N: 0.0001 to 0.010%, O:0.0001 to 0.020%, Cr: 0.001 to 0.50%, Mo: 0.001 to 0.10%, Nb: 0.001 to0.10%, V: 0.001 to 0.10%, Cu: 0.001 to 0.10%, W: 0.001 to 0.10%, Ta:0.001 to 0.10%, Ni: 0.001 to 0.10%, Sn: 0.001 to 0.050%, Sb: 0.001 to0.050%, As: 0.001 to 0.050%, Mg: 0.0001 to 0.050%, Ca: 0.001 to 0.050%,Y: 0.001 to 0.050%, Zr: 0.001 to 0.050%, La: 0.001 to 0.050%, and Ce:0.001 to 0.050%.
 3. A method for producing a steel sheet with excellentcold workability during forming according to claim 1, said methodcomprising: subjecting a steel strip having an ingredient compositionaccording to claim 1 to hot rolling by heating, followed by completingthe finish hot rolling at a temperature range of 800° C. or higher and900° C. or lower; coiling said hot-rolled steel sheet at a temperatureof 400° C. or higher and 550° C. or lower; pickling said hot-rolledsteel sheet, and then subjecting said hot-rolled steel sheet to atwo-step type annealing in which said hot-rolled steel sheet is retainedin two temperature ranges, wherein the two-step type annealing comprises(i) subjecting said hot-rolled steel sheet to a first step annealingperformed by retaining said hot-rolled steel at a temperature range of650° C. or higher and 720° C. or lower for 3 hours or longer and 60hours or shorter, and then a second step annealing performed byretaining the hot-rolled steel at a temperature range of 725° C. orhigher and 790° C. or lower for 3 hours or longer and 50 hours orshorter, and thereafter (ii) cooling said hot-rolled steel sheet to 650°C. or lower at a cooling rate of 1° C./hour or more and 30° C./hour orless.
 4. The method for producing a steel sheet according to claim 3,wherein the steel sheet has a cross-sectional shrinkage percentage of40% or more.
 5. A method for producing a steel sheet with excellent coldworkability during forming according to claim 2, said method comprising:subjecting a steel strip having an ingredient composition according toclaim 2 to hot rolling by heating, followed by completing the finish hotrolling at a temperature range of 800° C. or higher and 900° C. orlower; coiling said hot-rolled steel sheet at a temperature of 400° C.or higher and 550° C. or lower; pickling said hot-rolled steel sheet,and then subjecting said hot-rolled steel sheet to a two-step typeannealing in which said hot-rolled steel sheet is retained in twotemperature ranges, wherein the two-step type annealing comprises (i)subjecting said hot-rolled steel sheet to a first step annealingperformed by retaining said hot-rolled steel at a temperature range of650° C. or higher and 720° C. or lower for 3 hours or longer and 60hours or shorter, and then a second step annealing performed byretaining the hot-rolled steel at a temperature range of 725° C. orhigher and 790° C. or lower for 3 hours or longer and 50 hours orshorter, and thereafter (ii) cooling said hot-rolled steel sheet to 650°C. or lower at a cooling rate of 1° C./hour or more and 30° C./hour orless.
 6. The method for producing a steel sheet according to claim 5,wherein the steel sheet has a cross-sectional shrinkage percentage of40% or more.